Carbon/carbon composites are presently used for various nigh-temperature applications. Because such composites are dimensionally stable and retain their strength up to 3000.degree. C., they are an excellent material for areas where high temperature thermal protection is required.
The composites are used in applications such as the nose cap and leading edges of the shuttle orbiter where temperatures approach 3000.degree. C. during atmospheric re-entry and exceed even the temperature limits of the protective ceramic tiles used on the other surfaces of the orbiter. Carbon/carbon composites have a near zero coefficient of thermal expansion and a high heat capacity and are non-ablative materials that can withstand the thermal shock of re-entry, going from -160.degree. C. in the cold of space to 3000.degree. C. Unlike the ablative heat shields used in early manned spacecraft, carbon/carbon composites can undergo thermal cycling thousands of times with little decrease in properties. Also, carbon/carbon composites can withstand the impact of small meteors traveling at thousands of miles an hour without fracturing or developing cracks. Because of this stability and toughness, the carbon/carbon composites on the shuttle do not need to be replaced after every mission.
Carbon/carbon composites are also used in aircraft engines because of their high temperature performance. Nozzle components on fighter engines are constructed from carbon/carbon composites, not only because they are lighter and stronger than conventional materials, but also because they can withstand the temperatures and forces exerted on them by the engine exhaust gases. In fact, unlike the titanium super-alloys typically used in these applications, carbon/carbon composites actually increase in strength as the temperature increases. The turbine rotors in these engines often are often manufactured from carbon/carbon composites because they perform well at such high temperatures, thereby increasing the engine's efficiency.
Unfortunately, even though these materials are very tough, their strengths heretofore have normally been insufficient to allow them to be used as load carrying members. Also, because carbon/carbon composites are difficult to process and expensive, their use in industry has been limited to applications where high temperature performance, rather than cost, is the concern.
Known carbon/carbon composites are typically manufactured by infiltrating a woven carbon fiber preform with a polymeric resin. The impregnated preform is then heat treated in an inert atmosphere under pressure to produce a carbon fiber structure bound by a carbonaceous matrix. The initial carbon/carbon composite formed by this method is very porous, with a density less than water, which leads to dimensional instability and low mechanical properties. To alleviate these problems, the porous composite structure must be subsequently re-impregnated with the polymeric resin and further heat treated to densify the composite. This densification step is repeated until a density of about 1.6 grams per cubic centimeter is achieved to produce a composite with the desired mechanical properties and dimensional stability at elevated temperatures. These densification steps usually take several months, however, to perform and, thus, limit the use of the carbon/carbon composites to projects where cost and time are not an issue.
Carbon/carbon composites contain a primary carbon filler that is normally in fiber form and secured by a secondary carbon matrix binder. Some of the earliest commercial carbon/carbon composites utilized polycrystalline carbon as the filler and a coke as the binder. The polycrystalline carbon gave the composite its high mechanical properties. In today's advanced carbon/carbon composites, desirable components include a high modulus carbon fiber as the filler and a mesophase carbon as the binder. Such advanced composites have strength-to-weight ratios that are well above those of metals and ceramics.
Currently in the manufacture of many carbon/carbon composites, the carbon reinforcing fibers are first woven or braided into a three-dimensional (3-D) preform. This 3-D structure of carbon fibers increases the interlaminar shear strength of the final composite. Next, the 3-D preform is infiltrated with a polymeric or pitch resin that serves as the matrix of the composite. Polymeric resins typically have very low carbon yields, ranging from 40% to 50%. On the other hand, the carbon yield for an isotropic pitch can be as high as 70% to 80% and that for a mesophase pitch can reach 80% to 90%. After infiltration, the impregnated structure is pyrolyzed at temperatures ranging from 800.degree. C. to 1500.degree. C. to drive off all substances except the carbon. Unfortunately, the gaseous volatiles create pores and bloat the composite as they escape, reducing both the density and mechanical properties of the composite. To fill these pores and improve the mechanical properties of the composite, several densification cycles must be performed as mentioned above.
Densification consists of re-impregnating the porous composite with the resin and again pyrolyzing. This procedure yields a densified carbon/carbon composite, suitable for high temperature structural applications. The repeated cycles, however, can take several weeks for a single composite depending on the resin used. Heretofore, polymeric resins have been considered more suitable for impregnation due to their lower viscosity, but the low carbon yield makes many costly densification cycles necessary. Obviously, if the cost of carbon/carbon composites is to be reduced, new methods of manufacturing must be developed.
Recently, carbon fibers have replaced polycrystalline carbon as the preferred filler in carbon/carbon composites because of their high strength and elastic properties. Carbon also exhibits higher mechanical properties in the fiber form than in the original bulk material. With brittle materials, like carbon, this increase in properties is caused by both an increase in molecular orientation and a decrease in strength-limiting flaws. The decrease in flaw density due to the reduction of the cross-sectional area, combined with the increase in molecular orientation, makes the mechanical properties of carbon fiber approximately 500 times greater than that of bulk carbon.
Today, two different types of carbon fiber are commercially available: polymeric-based and pitch-based. The most common polymeric carbon fiber is produced from polyacrylonitrile (PAN). Numerous PAN-based carbon fibers have reactive functional groups at the surface that can react with the matrix of the composite. The reactive surface and the rough skin texture of polymeric-based carbon fibers create a strong fiber/matrix interface in the composite through chemical and physical bonding. The strong interface permits stress to be transferred from the matrix to the fiber, an important characteristic for fiber-reinforced plastics.
Pitch-based fibers, however, are the desired fibers for use in carbon/carbon composites. Mesophase pitch-based carbon fibers exhibit a more ordered graphitic structure than PAN-based carbon fibers. Such ordered structures provides fibers that are less susceptible to oxidation. They also exhibit a variety of transverse microstructures (morphology). Different fiber microstructures can be produced by varying the spinning conditions of the liquid crystalline precursor. Because the fibers are melt spun, mesophase pitch-based fibers have a smooth surface. During subsequent pyrolysis, the fiber loses only 10% to 20% of its mass, causing little collapsing of the fiber surface. Thus, the surface of the final pitch-based carbon fiber is much smoother than that of PAN-based fibers. Unlike PAN-based carbon fibers, the final mesophase pitch-based fiber has very few reactive functional groups on its surface. Because of the less reactive fiber surface and the optically smooth skin, composites made from pitch-based carbon fibers exhibit weaker interfacial bond strengths than those formed using PAN-based carbon fibers. The problems created by the weaker interfacial bond strengths may be overcome by well-known techniques of surface treating or sizing the pitch-based fibers.
Two types of matrix precursors may be used to make carbon/carbon composites: (1) those that form "hard" carbon residues, and (2) those that form "soft" carbon residues. The hard carbons are a non-graphitizing residue and are formed by pyrolysis of a infusible precursor, such as a thermoset polymer. Because the precursor is highly networked and crosslinked, it is unable to reorient and form graphite crystals during pyrolysis, regardless of the heat treatment temperature. Because hard carbons do not have the preferred orientation of graphite, composites formed using hard carbon matrices tend to exhibit lower Young's moduli. The crosslinking, however, does provide for stress transfer between the graphite planes, thus leading to increased toughness of the matrix. High toughness of the matrix and a moderate Young's modulus are typical for all the hard carbons, whereas a high modulus and a shear sensitivity of the matrix are characteristic of soft carbons.
Soft carbons are formed by the liquid-phase pyrolysis of polyaromatic resins, such as pitch. During pyrolysis, at a temperature of approximately 400.degree. C., aromatic polymerization converts any type of pitch matrix resin to a mesophase pitch, in which the larger polycarbon planes reorient themselves parallel to each other. These pre-oriented crystal-like regions of the liquid tend to coalesce, thus promoting crystallization and the formation of regions of perfect graphitic crystallinity upon further heat treatment. The adjoining layer planes within soft carbon are only weakly bound by van der Waals forces, making this material shear-sensitive. However, it is this weak interplanar bonding that allows the layer planes to reorient easily during pyrolysis.
Most raw pitches (Ashland 240, Ashland 260, etc.) are subjected to further processing prior to use as matrix resins in composites. These treatments increase the carbon yield of the pitch and, thus, reduce somewhat the number of densification cycles necessary during composite formation. Several grades of Aerocarb pitch are produced by the Ashland Petroleum Company by varying the heat treatment temperature and time. The product number (Aerocarb 60, 80, etc.) indicates the minimum coking value of the product.
When either polymeric or pitch matrix materials are used in conventional carbon/carbon composites, the resulting composites nearly always contain voids after pyrolysis. As previously mentioned, a major cause of porosity is insufficient wetting and/or inadequate infiltration of the precursor into the bundles of fibers in the preform. Usually the precursor is able to fill the large gaps found in a woven preform, but often it cannot infiltrate the small pores in the bundle, resulting in encapsulated regions of dry fibers. The resulting porosity is quite evident in green composites.
Another leading cause of porosity results from the differences in thermal behavior of the fibers and the matrix. Strong interfacial bonding in the green composite, combined with significant differences in the thermal expansion behavior of the matrix and fiber, can lead to fiber/matrix debonding during pyrolysis. The debonding creates stress concentrations during loading of the composite, resulting in premature cracking of the matrix and fracture of the fibers. Slow heating of the composite during carbonization somewhat alleviates, but does not eliminate, this problem.
The third and most significant cause of pore formation is mass loss during pyrolysis. As previously mentioned, during this heat treatment step the matrix precursor reacts and produces low molecular weight gaseous products, such as methane, ammonia, CO, and CO.sub.2. These pyrolysis gases cause bloating of the bulk matrix (similar to foaming) and, thus, create the well-known high porosities. Once formed, these pores can only be filled by redensification of the structure. The porosity and, thus, the need for redensification can be reduced by using a matrix precursor with a high carbon yield, thus reducing the amount of gases given off during pyrolysis.
Another technique to decrease bloating is through crosslinking or stabilization. If the matrix can be crosslinked prior to pyrolysis, bubbling and bloating can be minimized and increased carbon yields can be obtained. Even in crosslinked matrices, however, significant weight loss still occurs during pyrolysis, leading to shrinkage of the matrix and microcracking. The microcracking mechanisms typically observed in composites are delamination and fold-sharpening, characterized by shrinkage, distortions, and layer-rupture acting on the curved layer of the mesophase sheath around the fibers. This leads to fiber/matrix debonding and separation (delamination) of the graphitic layer planes in the matrix.
Carbon/carbon composites are normally formed from polymeric resins in the following general manner. First, a carbon fiber/polymeric composite is fabricated, either by wet-winding a one-dimensional (1-D) or two-dimensional (2-D) structure, or by melt-impregnating a multidimensional preform. Multidimensional preforms are preferred because they minimize the bulk matrix shrinkage during pyrolysis and improve the interlaminar shear strength of the final composite. Following impregnation, the preform is cured in an autoclave to form the green composite. The green composite is then carbonized in an inert atmosphere at temperatures ranging from 800.degree. C. to 1500.degree. C. to yield a hard carbon residue. During this step, the volatiles escape, thus creating a porous composite that, heretofore, have contained as much as 30% voids, by volume. The porous structure formed thereby is then densified by melt impregnation with additional resin, followed by subsequent recarbonization. As many as five densification cycles may be performed, requiring several weeks or months to achieve desired porosities of less than 5%.
Chemical vapor infiltration (CVI) is another technique commonly used to densify the porous structure. With CVI, a carbon-rich atmosphere is created by cracking hydrocarbon gases at temperatures ranging from 1000.degree. C. to 2000.degree. C. The preform is placed in this atmosphere and maintained at a slightly lower temperature than the surrounding atmosphere which causes the carbon to precipitate onto the preform and create a carbon matrix. This method is extremely expensive, time consuming, and difficult to control.
When pitch resins are employed, the matrix must be oxidized in order to prevent bloating as the pyrolysis gases evolve during carbonization. With mesophase pitch, the oxidation also serves to preserve the orientation of the crystallites while in the liquid crystalline state. In order to effectively crosslink the carbon layers, the pitch is oxidized for many hours (50 to 100) at temperatures slightly below the softening point. The oxidized composite is then carbonized at temperatures from 800.degree. C. to 1500.degree. C. in an inert atmosphere. Even though volatiles and pyrolysis gases still can create some porosity, the oxidation step lowers the porosity levels to between 10% and 20%. In such prior art processes, the composite is densified to the desired level by repeated melt impregnations or CVIs.
High carbon-yield pitches usually have higher viscosities than the phenolic resins used for carbon/carbon composites. Such high viscosity, however, makes it difficult to fill pitch into the fine pores within the bundles of fibers, resulting in a slightly higher porosity. However, because composites formed using high carbon-yield pitches require fewer redensification cycles, this increase in porosity has previously been tolerable.
Even though impregnation can be difficult, pitch resins easily wet carbon fibers and provide good adhesion. However, the difference in the thermal expansion behavior of the fiber and the pitch makes strong bonding at the fiber/matrix interface detrimental to the formation of the composite during repeated densification. Pitch has a positive coefficient of thermal expansion, while the fiber has a negative coefficient of thermal expansion. Thus, stresses develop and remain in the composite after processing. If interfacial bonding is strong, the stresses will be relieved by matrix cracking and/or fiber failure. Debonding also will occur, increasing the porosity of the composite. The smooth, unreactive surfaces of mesophase pitch-based carbon fibers, however, reduce adhesion. As a result, mesophase pitch-based carbon fibers are used in most advanced carbon/carbon composites. Many of these problems would be alleviated if a composite forming process that eliminates the repeated densification cycles was available.
At present, the high cost of carbon/carbon composites precludes their use in high-volume applications. Although carbon/carbon composites are more suitable for many applications, lower cost materials, such as ceramics, are selected. In order for carbon/carbon composites to compete with ceramics and enter high-volume markets, their manufacturing costs must be drastically reduced. A need exists for a technique to reduce the number of densification steps and thereby reduce the cost of carbon/carbon composites. The present technique overcomes these and other shortcomings of the prior art by utilizing a continuous powder coating process to produce a pitch-based towpreg that can be fabricated readily into thick-walled woven carbon/carbon structures.